**3. Results**

The GDOES data (Table 1) were used to quantify the thickness of the coatings and assess the homogeneity of the elemental distribution within the coatings. The elemental profiles show all the main elements that make up the target, such as tantalum, zirconium, silicon, and boron. In coatings applied in environments containing nitrogen and ethylene, there is also a signal from nitrogen and carbon, respectively (Figure 1). The emergence of Ti lines signifies the boundaries between the coatings and Ti-alloys substrate and allows for an assessment of the coatings' thickness (6.6 μm in the case of coating 3 and 5.8 μm in the case of coating 5).

**Figure 1.** Typical glow discharge optical emission spectroscopy (GDOES) profiles of coatings 3 (10 sccm N2) (**a**) and 5 (10 sccm C2H4) (**b**).

According to the GDOES data, all elements were evenly distributed across the thickness of coatings. The concentration of oxygen impurities was below 1 at.%. The minor oxygen contamination resulted from the presence of corresponding impurities both in the ceramic target and in the working gas. The depth-averaged coating composition is shown in Table 1.

The coatings deposited in argon contained 47.6 at.% of metals and 52.4 at.% of non-metals, so the metal-to-nonmetal species ratio was close to 1. The increased flow of reactive gases resulted in higher nitrogen or carbon content in coatings and a decreased concentration of elements inherited from the target (except for Zr). The maximum N and C concentrations (42.4 and 30.3 at.%, respectively) were achieved at 10 sccm of N2 and C2H4, correspondingly.

Non-reactive deposited coating 1 had a dense structure with pronounced columnar grains, which is typical for ion-plasma deposited coatings [34,35] (Figure 2a).

**Figure 2.** Cross-section scanning electron microscopy (SEM) images of as-deposited coatings 1 (Ar) (**a**), 2 (5 sccm N2) (**b**), 3 (10 sccm N2) (**c**), 4 (5 sccm C2H4) (**d**), and 5 (10 sccm C2H4) (**e**).

A similar columnar structure was reported for TaSi2-based coatings in [18,28]. When N2 (coatings 2 and 3) or C2H4 (coatings 4 and 5) were added to the working gas, the columnar structure became less pronounced or completely disappeared. The coatings produced by reactive sputtering featured a nearly perfect, defect-free structure.

The coatings' thickness, depending on the deposition modes, ranged from 6.3 to 8.0 microns (Table 1). The introduction of a minor amount of nitrogen (coating 2) had no significant effects on the growth rate (180 nm/min). With an increase in nitrogen concentration (coating 3), a boost in the growth rate of up to 200 nm/min was observed. When carbon-containing gas was added to the sputtering atmosphere, the growth rate decreased to 157 nm/min (coating 4) or 175 nm/min (coating 5).

The XRD patterns and HRTEM-imaged structure of the coatings 1–5 are shown in Figure 3.

**Figure 3.** X-ray diffraction analysis (XRD) patterns of coatings 1(Ar), 2 (5 sccm N2), 3 (10 sccm N2), 4 (5 sccm C2H4), and 5 (10 sccm C2H4) deposited onto alumina substrate (**a**); selected area electron diffraction (SAED) and high-resolution transmission electron microscope (HR TEM) image of coatings 3 (10 sccm N2) (**b**) and 5 (10 sccm C2H4) (**c**).

A signal from the substrate of Al2O3 (JCPDS 88-0107 card) was detected for all samples. The XRD pattern for coating 1 contained peaks associated with reflections from planes (100), (102), (111), (210), (203), (220), and (115) of the h-TaSi2 hexagonal phase (JCPDS 89-2941), similar to [36]. The size of the h-TaSi2 crystallites determined by the Scherrer formula for the coating obtained in Ar was approximately 11 nm. The transition to reactive sputtering (in Ar-N2 and Ar-C2H4) resulted in a considerable structural refinement of the coatings. The most intense peak in the 2Θ = 25–45◦ range for coatings obtained in Ar-N2 can be attributed to the Ta-N bonds in the coating (FCC-TaN, JCPDS 89-5198), which agrees with the previously obtained XRD results for Ta-Si-N coatings deposited at a different N2/(N2 + Ar) ratio [37]. The crystallite size of the phase h-TaSi2 for coatings 2 and 3 produced by reactive sputtering was evaluated according to the most distinct lines. Coating 2, deposited at a flow rate of 5 sccm N2, had a grain size of the h-TaSi2 phase equal to ~6 nm. As the N2 flow rate increased, the crystallite size decreased to 4.5 nm. Coatings 4 and 5 deposited in the Ar + C2H4 atmosphere show broad peaks on XRD patterns at 2Θ from 30◦ to 45◦. Their position can be explained by the presence of Ta-Si and Ta-C bonds (TaC, JCPDS 89-3831). The introduction of nitrogen and carbon into the coatings resulted in grain refinement of h-TaSi2 phase to 3–6 nm according to HR TEM иselected area electron diffraction (SAED) data (Figure 3b,c) The sizes of the h-TaSi2 phase crystallites for coatings deposited at 5 and 10 sccm C2H4 flow rates were equal to 3.5 and 3.0 nm, respectively. The TEM data correlated well with the results of XRD.

The mechanical properties of coatings, such as hardness (H), elastic modulus (E), elastic recovery (W), plasticity index (H/E), and the resistance to plastic deformation (H3/E2) and compressive stress (σ), are shown in Table 2.


**Table 2.** Mechanical and chemical properties of coatings.

Coating 1 deposited in Ar demonstrated H = 12.5 GPa, E = 208 GPa, and W = 43.4%. The transition to reactive deposition at a flow rate of 5 sccm N2 resulted in an increase in H by 46% as well as an increase in E and W by 12% and 32%, respectively. When the N2 flow rate was increased to 10 sccm, H rose further to 29.2 GPa, while E and W increased to 279 GPa and 77.9%, respectively. This effect might be related to the formation of tantalum nitride phase [38,39] or structure refinement [40]. The introduction of carbon atoms into the coatings also resulted in an increase of mechanical properties. Coating 4 deposited at a flow rate of 5 sccm C2H4 had H = 21.3 GPa, E = 267 GPa, and W = 62.4%, which is 70%, 28%, and 44% higher, respectively, than the values obtained for the non-reactive sample. Coating 5 demonstrated hardness H = 28.3 GPa, Young's modulus E = 288 GPa, and elastic recovery W = 76.4%. The mechanical properties of coatings increased at higher carbon concentrations, similar to the results obtained in [41,42]. The increase in mechanical properties can be partially attributed to the increase of compressive stresses in the carbon-rich coatings (Table 2). Table 1 shows the values of the parameters H/E and H3/E2, which are important in terms of assessing the level of wear resistance and determining the mechanism of localized deformation [43,44]. In indentation experiments, the coatings with H3/E2 below 0.5 GPa are deformed in a heterogeneous fashion with the associated formation of shear bands, whereas the coatings with H3/E<sup>2</sup> above 0.6 GPa the deformation are homogeneous, with no structural transformations of the surface [44].

The tribological testing of the coatings was carried using counterparts made of 100Cr6 and Al2O3, which differ significantly from each other in terms of physical and mechanical characteristics. The results of tests determining the coatings' friction coefficient are shown in Figure 4.

**Figure 4.** Friction coefficient vs. distance dependence during testing of coatings against 100Cr6 (**a**) and Al2O3 (**b**) counterparts. Inserts show the values of wear rate for coatings and counter-bodies.

When the 100Cr6 ball was used as a counterpart, the behavior of coatings 1–3 was similar. The friction coefficient increased from the initial values of 0.2–0.3 to 0.9–1.1 at a distance of 10–20 m and then stabilized at 1.04 (coating 1), 0.98 (coating 2), and 0.92 (coating 3) and remained relatively constant until the end of the measurements at a distance of 50 m. The behavior of coating 4 obtained at 5 sccm C2H4 at the initial moment of the test was similar to that of coatings 1–3. However, when reaching a distance of 17.5 m, the friction machine automatically stopped the measurement as a result of exceeding the permissible friction coefficient (>1.1). Coating 5 with the highest carbon content showed the lowest and the most stable *f* ~0.3 at a distance up to 25.7 m, followed by a jump to 0.65 (25.8–32.4 m), and further monotonous increase from 0.39 up to 0.76 at the distance from 32 to 50 m. This behavior can be explained by the positive role of carbon [45,46], as well as by specific tribochemical reactions [47].

In the tests where Al2O3 ball was used as a counterpart, for coatings 1–3, the friction coefficient *f* increased sharply to values of 0.95–1.05 at a 0–2 m distance. This resulted in the premature termination of tests for samples 1 and 3 at a distance below 10 m as a result of exceeding the permissible values of *f*. The friction coefficient of sample 2 obtained at 5 sccm N2 was stable after a short run-in period and remained within a 0.95–1 range until the end of the test at 50 m. Coating 4 deposited in a mixture of Ar and C2H4 showed a rapid rise of the f value to 0.8 and then remained stable until the end of the test. Coating 5 with a high carbon content had similar behavior with slightly lower values at distances up to 20 m.

Figure 5 provides three-dimensional profiles of wear tracks after testing with 100Cr6 and Al2O3 balls for coatings 1, 2, and 4.

**Figure 5.** 3D images of wear tracks on the surface of coatings 1(Ar), 2 (5 sccm N2), and 5 (10 sccm C2H4) after testing against 100Cr6 (**a**–**c**) and Al2O3 (**d**–**f**).

When tested against the steel ball counterpart, none of the coatings 1–5 showed any signs of surface wear; however, noticeable pile-ups of wear debris were present, suggesting that wear of steel ball and adhesion of steel debris to the coatings were the main wear mechanisms. Coatings 1, 2, and 4 had pronounced areas of wear debris accumulation, concentrated in local zones and having a height of 3–5 microns. In the case of coatings 3 and 5 with the maximum nitrogen and carbon concentration, the wear debris was distributed more evenly in the tribocontact zone and had a height of 2–3 microns. The wear rate (Vw) of the steel ball used to test coating 1 was 7.9 <sup>×</sup> 10−<sup>5</sup> mm3·N−1·m−<sup>1</sup> (Figure 4a, insert). The highest (2.0 <sup>×</sup> 10−<sup>4</sup> mm3·N−1·m<sup>−</sup>1) and lowest (8.4 <sup>×</sup> 10−<sup>6</sup> mm3·N−1·m<sup>−</sup>1) ball wear rates were achieved for nitrogen-rich coatings 2 and 3. The increase of carbon content was associated with the decrease of Vw from 5.9 <sup>×</sup> <sup>10</sup>−<sup>5</sup> (coating 4) to 2.5 <sup>×</sup> <sup>10</sup>−<sup>5</sup> mm3·N−1·m−<sup>1</sup> (coating 5).

During the testing with alumina ball counterpart, coating 1 deposited in Ar experienced uneven wear; that is, both segments with no visible wear and areas of complete wear with a depth of ~8 microns were present (Figure 5d). The low wear resistance of coating 1 can be associated with its pronounced columnar structure, characterized by low fracture toughness [35]. The wear rate of the coating was 5.4 <sup>×</sup> 10−<sup>3</sup> mm3·N−1·m−<sup>1</sup> (Figure 4b, insert). In coatings 2–5, deposited by reactive sputtering, the wear depth did not exceed the coating thickness. In the case of nitrogen-containing coating 2, wear track with a width of 450 microns and depth below 2 microns was observed. The wear rate of the coating was 0.3 <sup>×</sup> 10−<sup>3</sup> mm3·N−1·m−1, which is ~20 times lower than the values obtained for the coating 1 deposited in Ar. Coating 3 with the maximum nitrogen content showed no signs of wear (test run ~2 m). Sample 4 obtained at 5 sccm C2H4 demonstrated a wear rate of 0.4 <sup>×</sup> <sup>10</sup>−<sup>3</sup> mm3 <sup>N</sup>−<sup>1</sup> <sup>m</sup><sup>−</sup>1, which

is equivalent to the value for a coating deposited in nitrogen at the same gas flow rate. When the C2H4 flow rate increased to 10 sccm, an increase in the wear rate was observed (1.3 <sup>×</sup> 10−<sup>3</sup> mm3·N−1·m<sup>−</sup>1). Our previous work [23] addressed the case of Ta-Si-C and Ta-Si-C-N coatings produced by sputtering of composite SHS-cathodes TaSiC. The nitrogen-rich coatings (Ta-Si-C-N) showed a low wear rate against Al2O3 ball (1.7–2.8 <sup>×</sup> 10−<sup>5</sup> mm3·N−1·m−1, whereas their nitrogen-free counterparts suffered complete wear.

The wear areas on alumina balls show that the lowest achieved wear rate was 1.6 <sup>×</sup> 10−<sup>5</sup> mm3·N−1·m−<sup>1</sup> and was achieved for coating 1 (Figure 4b, insert). Nitrogen-rich coatings 2 and 3 were worn at rates 3.3 <sup>×</sup> 10−<sup>4</sup> and 4.8 <sup>×</sup> 10−<sup>5</sup> mm3·N−1·m−1, respectively. The wear of counter-bodies for coatings 4 and 5 was 2.2 <sup>×</sup> 10−<sup>4</sup> and 5.9 <sup>×</sup> 10−<sup>5</sup> mm3·N−1·m<sup>−</sup>1, respectively. Therefore, the increase in both carbon and nitrogen content in reactively-deposited coatings is beneficial for the decreased wear of alumina counter-bodies.

The SEM images of the surface of coatings 1–5 annealed at the maximal temperature of 1000 ◦C are shown in Figure 6.

**Figure 6.** Top-view SEM images of coatings 1(Ar) (**a**,**b**), 2 (5 sccm N2) (**c**), 3 (10 sccm N2) (**d**), 4 (5 sccm C2H4) (**e**), and 5 (10 sccm C2H4) (**f**) after air annealing with maximal temperature of 1000 ◦C. Insert on 1a shows the energy dispersive spectroscopy (EDS) map for Ta signal.

The surface of coating 1 deposited in Ar contained visible cracks and delaminations. However, a local uniform microstructure with no visible defects can be seen at ×2000 magnification. The cracks formation likely resulted from the difference in thermal expansion coefficients between the coating and the substrate. As a result, tensile stresses develop along the perimeter of the defect, which leads to the coating's destruction [48]. For coating 2 obtained at 5 sccm N2, a specific structure (Figure 6b), resembling blisters or bubbles was observed after the annealing. The formation of cavities or pores can be associated with the oxidation of nitrogen-containing phases in the coating and subsequent release of gaseous nitrogen oxides through the oxide film at high temperatures [37]. The bubbles occupied about 85% of the surface area of coating 2. When the N2 flow rate increased to 10 sccm (coating 3), a reduction in the number of bubbles was observed and the total area of defects was ~10% of the surface area. Samples 4 and 5 showed the same defective structure, with open pores (~0.1–2 microns) occupying up to ~75% of the surface. The formation of pores is caused by the accumulation of COx in the upper layer of the coating and the formation of blisters, which, under a certain gas pressure, burst and form pores on the surface [49].

SEM cross-section images of annealed coatings 1–5 are shown in Figure 7.

**Figure 7.** Cross-section SEM images of coatings 1 (Ar) (**a**,**b**), 2 (5 sccm N2) (**c**), 3 (10 sccm N2) (**d**), 4 (5 sccm C2H4) (**e**), and 5 (10 sccm C2H4) (**f**) after air annealing with maximal temperature of 1000 ◦C.

Coating 1 has experienced delamination from the substrate during the annealing (Figure 7a) and the formation of oxide layers on both sides of the coating. The oxide layer was pore-free, with a maximum thickness of 3.5 microns, and was comprised mainly of SiOx and TaOx. Cross-section SEM images of sample 2 show a porous oxide layer with a thickness of 4.4 microns. Pores (d = 0.2–2.0 microns) were concentrated in the upper part of the oxide layer. With an increase in the nitrogen content, the thickness of the oxide layer decreased by 23% and amounted to 3.4 microns. Coating 4 obtained at 5 sccm C2H4 showed an oxide layer thickness of 3.6 microns. The formed SiOx-TaOx oxide film had a loose structure with a pore size up to ~2.6 microns. The oxide layer of sample 5 with a thickness of 3.8 microns also had a porous structure. Pores with 0.2–1.5 microns diameter were randomly distributed over the thickness of the oxide layer. Coating 2 obtained at 5 sccm N2 had the maximum thickness of the oxide layer (Table 2), and the minimum thickness of the oxide layer was characteristic of sample 3, deposited at a flow rate of N2 10 sccm.
